Hot Rolled Precipitation Strengthened and Grain Refined High Strength Dual Phase Steel Sheet Possessing 600 MPa Minimum Tensile Strength and a Process Thereof

ABSTRACT

A process for producing dual phase steel sheet including steps of making a liquid steel having a chemical composition in wt % of C: 0.03-0.12. Mn: 0.8-1.5. Si: &lt;0.1, Cr: 0.3-0.7, S: 0.008 maximum, P: 0.025 maximum, Al: 0.01 to 0.1, N: 0.007 maximum. Nb: 0.005-0.035. and V: 0.06 maximum, remainder Fe; continuous casting the liquid steel into a slab; hot rolling the slab into a hot rolled sheet at finish rolling temperature (FRT) 840±30 ° C.; cooling the hot rolled sheet on the run out table at a cooling rate 40 −70° C./s to an intermediate temperature (Tint) of 720° C.≤Tint≤650° C.; natural cooling the hot rolled sheet for a duration of 5-7 seconds and rapidly cooling the hot rolled sheet to transform remaining carbon enriched austenite to martensite, at cooling rate of 40-70 ° C./s to a coiling temperature below 400° C.

CROSS-REFERENCE TO RELATED APPLICATIONS

This application is the United States national phase of International Application No. PCT/IN2017/050171 filed May 10, 2017, and claims priority to Indian Patent Application No. 201731004831 filed Feb. 10, 2017, the disclosures of which are hereby incorporated in their entirety by reference.

BACKGROUND OF THE INVENTION Field of the Invention

The present disclosure relates to a process for producing hot rolled high strength dual phase steel. The disclosure further relates to hot rolled high strength dual phase steel with >600 MPa tensile strength and 25% total elongation.

Description of Related Art

Motor vehicle fuel consumption and resultant emission is one of the major contributors to air pollution. Light-weight environmental friendly vehicle design is required to address the problems of environmental pollution. Successful light-weight motor vehicles require utilization of advanced high strength high strength steel (AHSS) sheets. However, because of its poor formability, the AHSS sheet cannot be applied easily to a wide variety of motor vehicle components. Hence, the ductility and formability required for AHSS sheet becomes increasingly demanding. Therefore addressing the present scenario has necessitated development of a hot rolled steel sheet with high tensile strength coupled with excellent uniform elongation, working hardening rate and total elongation for automotive component such as wheel web applications.

Hence, in order to replace the existing grades of steel used for automotive structural and wheel web applications, it is necessary to develop hot-rolled steel sheets which not only possess a minimum tensile strength of 600 MPa but also have good formability and good surface quality.

European patent EP1398392A1 and U.S. Pat. No. 8,337,643 disclose a method of producing a hot rolled dual phase (ferrite+martensite) steel of minimum tensile strength of 590 MPa. Though the proposed steels achieved the strength level, it contains high amount of Si (minimum 0.5 wt. % in European patent and 0.2 wt. % in US patent). Presence of Si will lead to surface scales, generally called as tiger marks.

European Patent EP2053139B1 discloses a method in which a hot rolled steel sheet is subjected to heat treatment after forming so as to achieve a tensile strength varying in the range of 440 to 640 MPa.

However, the heat treatment after forming, which is an essential part of the disclosure, is likely to add to the processing cost and hence is not suitable for mass production.

European Patent EP2578714A1 discloses a method of producing hot-rolled steel sheets with a minimum tensile strength of 590 MPa with excellent bake hardenability and stretch-flangeability. According to the proposed method the steel must contain 1.7 to 2.5 wt % of Mn. When added in such large amounts, Mn tends to segregate in the central portion in the thickness direction, which not only induces cracking during press forming but also leads to inconsistency in achieving the desired stretch-flangeability.

It is also important to understand automotive wheel to develop the steel. The automotive wheel is composed of a disk and a rim. While the disc is press formed, the rim is flared and then roll formed after flash butt welding. Therefore, the material needed to form the disk needs to have good deep drawability, stretch formability and stretchability, whereas the material needed to form the rim needs to have good formability after welding. After the wheel discs and rims are formed by their respective processes, they are assembled by means of spot welding or arc welding. Hence the materials for both rim and disc use need to have good spot weldability. From the point of view of application, the most important functional requirement for auto-wheels is durability, which can be increased by increasing the fatigue strength of the wheel material.

The various studies conducted in the recent past show that precipitation hardened steels and dual phase (DP) steels are both suitable for wheel disc application. From the fatigue strength considerations, the upper limit of the tensile strength of steels for wheel use is 600 MPa (or 85 ksi) [T, Irie, K. Tsunoyama, M. Shinozaki and T. Kato: SAE Paper No. 880695, 1988]. This is because when the tensile strength is increased beyond 600 MPa, the consequent increased notch sensitivity results in lowering of the fatigue strength. Hot rolled DP steels with a tensile strength of 600 MPa (or HR-DP 600) have become a very popular choice for wheel disc applications owing to their superior strength and formability and at the same time good stretchability (high n value) and spot weldability. However, it is difficult to produce the HR-DP 600 steel in any mill because many process parameters, e.g. the finish rolling temperature, cooling rate etc. are needed to be optimized and fine tuned keeping in mind the mill configuration e.g. the length of the run out table, water volume available etc. in order to obtain the desired microstructural features which in turn will decide the final mechanical properties. All the existing patents and literature have considerable amount of Si to increase ferrite strength in order to fatigue life of the steel.

SUMMARY OF THE INVENTION

In view of the foregoing limitations inherent in the prior-art, it is an object of the disclosure to propose process for producing hot rolled precipitation strengthened high strength dual phase steel sheet the tensile strength more than 600 MPa with lower percentage of Si.

Another object of the disclosure is to propose process of producing hot rolled precipitation strengthened high strength dual phase steel sheet, with lower percentage of Si.

Another object of the disclosure to propose hot rolled precipitation strengthened high strength dual phase steel sheet the tensile strength more than 600 MPa with lower percentage of Si.

Still another object of the disclosure is to propose hot rolled precipitation strengthened high strength dual phase steel sheet, with lower percentage of Si.

The disclosure provides a process for producing dual phase steel sheet. The process comprises steps of making a liquid steel having chemical composition in wt % of C: 0.03-0.12, Mn: 0.8-1.5, Si: <0.1, Cr: 0.3-0.7,S- 0.008 max, P-0.025 max, Al- 0.01 to 0.1, N- 0.007 max, Nb: 0.005-0.035, and V- 0.06 max; continuous casting the liquid steel into a slab; hot rolling the slab into a hot rolled sheet at finish rolling temperature (FRT) 840±30 deg. C; cooling the hot rolled sheet on Run Out Table at cooling rate 40-70° C./s achieving intermediate temperature (TINT) 720≤TINT≤650; natural cooling the hot rolled sheet for duration 5-7 seconds; and rapidly cooling the hot rolled sheet to transform remaining carbon enriched austenite to martensite, at cooling rate of 40-70 deg. C/s to achieve coiling temperature below 400 deg. C.

BRIEF DESCRIPTION OF THE ACCOMPANYING DRAWINGS

FIG. 1 illustrates various steps for a process for making a high strength dual phase steel in accordance with an embodiment of the disclosure.

FIG. 2 illustrates a schematic diagram of cooling profile to obtain the high strength dual phase steel in accordance with an embodiment of the disclosure.

FIG. 3 illustrates a tensile stress strain plot of strip 1 accordance with an embodiment of the disclosure.

FIG. 4 illustrates an optical micrograph of Strip 1 (Nital etched) in accordance with an embodiment of the disclosure.

FIG. 5 illustrates an optical image of Lepera etched sample white phase: Martensite (a); Dark phase: Ferrite (a′) in accordance with an embodiment of the disclosure.

FIG. 6 illustrates an optical image of Le pera etched sample: Fine grains as small as 2 μm can be noticed in accordance with an embodiment of the disclosure.

FIG. 7 illustrates scanning electron microscopy images of strip 1 in accordance with an embodiment of the disclosure.

FIG. 8(a) illustrates bright field TEM micrograph of one of the precipitates in the ferrite matrix; 8(b) Dark field image of 8(a); 8(c) selected area diffraction pattern from Nb(C,N) precipitate, 8(d) Dark field image showing Nb(C,N) precipitates; 8(e) EDS spectrum of the precipitate and 8(f) Composition of the precipitate

DESCRIPTION OF THE INVENTION

Various embodiments of the disclosure provide a process for producing dual phase steel sheet, comprising steps of making a liquid steel having chemical composition in wt % of C: 0.03-0.12, Mn: 0.8-1.5, Si: <0.1, Cr: 0.3-0.7, S- 0.008 max, P- 0.025 max, Al- 0.01 to 0.1, N- 0.007 max, Nb: 0.005-0.035, and V- 0.06 max;

continuous casting the liquid steel into a slab; hot rolling the slab into a hot rolled sheet at finish rolling temperature (FRT) 840±30 deg. C; cooling the hot rolled sheet on Run Out Table at cooling rate 40-70° C./s achieving intermediate temperature (TINT) 720 TINT 650; natural cooling the hot rolled sheet for duration 5-7 seconds; and rapidly cooling the hot rolled sheet to transform remaining carbon enriched austenite to martensite, at cooling rate of 40-70 deg. C/s to achieve coiling temperature below 400 deg. C.

Another embodiment of the disclosure provide A dual phase steel sheet, comprising a chemical composition in wt % C: 0.03-0.12, Mn: 0.8-1.5, Si: <0.1, Cr: 0.3-0.7, S- 0.008 max, P- 0.025 max, Al- 0.01 to 0.1, N- 0.007 max, Nb: 0.005-0.035, and V- 0.06 max.

Shown in FIG. 1 is a process (100) for producing dual phase steel sheet. At Step (104) a liquid steel is made. Following is the composition of the liquid steel (in wt. %) C: 0.03-0.12, Mn: 0.8 1.5, Si: <0.1, Cr: 0.3-0.7, S- 0.008 max, P- 0.025 max, Al- 0.01 to 0.1, N- 0.007 max, Nb: 0.005-0.035, and V-0.06 max.

The addition of each alloying element and the limitations imposed on each element are essential for achieving the target microstructure and properties.

C: 0.03-0.12%: Carbon is one of the most effective and economical strengthening elements. Carbon combines with Nb or V to form carbides or carbonitrides which bring about precipitation strengthening. This requires a minimum of 0.03% C in the steel. However, in order to have good weld-ability, the carbon content has to be restricted to less than 0.12%.

Mn: 0.8-1.5%: Manganese not only imparts solid solution strengthening to the ferrite but it also lowers the austenite to ferrite transformation temperature thereby refining the ferrite grain size. However, the Mn level cannot be increased to beyond 1.5% as at such high levels it enhances centerline segregation during continuous casting.

Si <0.1 wt. %: Silicon like Mn is a very efficient solid solution strengthening element. However, Si leads surface scale problems in hot rolling and hence it should be restricted to less than 0.1% in order to prevent the formation of surface scales.

Nb: 0.035% maximum: Niobium is the most potent microalloying element for grain refinement even when it is added in very small amounts. When in solid solution it lowers the austenite to ferrite transformation temperature which not only refines the ferrite grain size but also promotes the formation of lower transformation products like bainite. However, to ensure the effectiveness of Nb, it should not be allowed to precipitate before the transformation temperature is reached. To ensure that the entire Nb content remains in solution before rolling commences and it is alone added, the maximum Nb content is restricted to 0.035%.

V: 0.06% maximum: Microalloying by Vanadium also leads to precipitation strengthening as well as grain refinement. The solubility of Vanadium in austenite is more than that of other microalloying elements and so it is more likely to remain in solution prior to transformation. During phase transformation, vanadium precipitates as carbides and/or nitrides, depending on the relative carbon and nitrogen contents, at grain boundaries resulting in precipitation strengthening as well as grain refinement. In order to achieve the desired strengthening, it is required to add either Nb or V. Both can also be added. If V alone is added, it is required up to 0.06 wt. %.

P: 0.025% maximum: Phosphorus content should be restricted to 0.025% maximum as higher phosphorus levels can lead to reduction in toughness and weldability due to segregation of P into grain boundaries.

S: 0.008% maximum: The Sulphur content has to be limited otherwise it results in a very high inclusion level that deteriorates formability.

N<0.007: Too high N content raises the dissolution temperature of Nb(C, N) and hence reduces the effectiveness of Nb. Reducing nitrogen levels also positively affects ageing stability and toughness in the heat-affected zone of the weld seam, as well as resistance to inter-crystalline stress-corrosion cracking. Thus N levels should be preferably kept below 0.007.

Al 0.01 to 0.1: Al is used to remove undesirable oxygen from molten steel and hence steel contains some amount of Al, may be upto 0.05 wt. %. Excess (high) Al in steel making is a major problem as it decreases hot deformation of cast slab besides nozzle clogging during casting. Therefore, Al needs to be restricted to 0.1 wt. %.

At step (108) the liquid steel is continuously casted into a slab.

The liquid steel of the specified composition is first continuously casted either in a conventional continuous caster or a thin slab caster. When cast in a thin slab caster, the temperature of the cast slab is not allowed to drop to a temperature below 950° C. This is because if the thin slab temperature falls below 950° C., Nb precipitation occurs. Then it becomes difficult to completely dissolve the precipitates in the subsequent reheating process rendering them ineffective for precipitation strengthening.

Reheating: After casting the slab with the specified composition, the slab is reheated to a temperature of 1100 to 1200° C. for a duration of 20 minutes to 2 hours. The reheating temperature should be above 1100° C., to ensure complete dissolution of any precipitates of Nb or/and V that may have formed in the preceding processing steps. A reheating temperature greater than 1200° C. is also not desirable because it leads to grain coarsening of austenite and/or excessive scale loss.

At Step (112) the slab is hot rolled into a hot rolled sheet at finish rolling temperature (FRT) 840 ±30° C.

The hot rolling constitute of a roughing step above the recrystallization temperature and a finishing step below the recrystallization temperature, when rolling is done in a conventional hot strip mill. In case Continuous Strip Processing is used for producing this steel, where there is no separate roughing mill, the deformation schedule is designed in such a manner that the cast structure is destroyed in the initial stands and finishing is done below the recrystallization temperature. More specifically the finish rolling in either set up should be done at a temperature, T_(FRT) given by 840 +/−30° C.

Laminar cooling on the Run-Out-Table (ROT): At step (116) the hot rolled sheet is cooled on a Run Out Table at cooling rate 40-70° C./s. The said cooling rate is maintained to achieve intermediate temperature (TINT) 720 TINT 650.

The cooling rate should be higher than 4CPC/s to prevent formation of pearlite. Any pearlite, or degenerate pearlite if formed leads to deterioration of both, tensile strength as well as stretch flangeability. High cooling rate also results in lowering the ferrite start temperature which leads to refinement of the ferrite grain size. It also prevents the growth of the ferrite. By increasing the cooling rate and controlling rolling schedule, the desired grain size of 2-6 μm can be achieved. The cooling rate may not be more than 70° C./s because then the desired amount of ferrite will not form. This fast cooling is continued up to an intermediate temperature. The intermediate temperature (TINT) should be 650<TINT<720° C.

At step (120), the strip is allowed to naturally cool while being transferred on RoT. The duration of air cooling is critical and is 5 to 7 seconds. If the strip is allowed to cool for less than 5 seconds, then sufficient amount of ferrite will not be formed. On the other hand, if the strip is allowed to air cool more than 7 seconds then it will results in insufficient amount of martensite.

During this period, austenite transforms to ferrite. However, entire austenite will not transform to ferrite as time is not sufficient for complete transformation. As a result remaining austenite at the end of natural cooling will be enriched with carbon because ferrite cannot accommodate average carbon content in the steel.

At step (124) the strip is further cooled rapidly after naturally being cooled at step (120). This ensures the transformation of remaining carbon enriched austenite to martensite. The cooling rate during this period is 40-70° C./s to achieve coiling temperature below 400° C. The coiling temperature can be as low as 100 deg.C

The strengthening contributions from solid solution elements and microalloying elements are restricted. Also, the extent of possible grain refinement, by controlled rolling and cooling is limited to 2 μm due to which the high strength dual phase steel obtained.

The microstructure obtained comprises martensite particle/phase in the ferrite matrix. The microstructure is uniform or in other words martensite phase is distributed uniformly throughout the ferrite matrix. Furthermore, bainite or degenerate pearlite/pearlite and grain boundary cementite is avoided and high strength dual phase steel sheet achieves good work hardening rate, low yield point and continuous yielding. The contribution of each of the microstructural components is described below:

a) Ferrite: The hot rolled steel sheet according to the present disclosure has 75-90% ferrite (by vol.). The ferrite is strengthened by solid solution strengthening contributions from Mn. Using suitable processing conditions, the grain size is restricted to 2-5 μm. This grain refinement of ferrite leads to strengthening of the ferrite, the amount of which is decided by the Hall-Petch relationship. Also it is precipitation strengthened by the formation of fine Nb,V(CN) precipitates.

b) Martensite: The amount of martensite in the microstructure is 10-25% (by vol.). The strengthening from martensite comes from its structure, carbon content and higher dislocation density.

c) Bainite: The amount of martensite in the microstructure is less than 5% (by vol.).

The high strength dual phase steel sheet has got improved fatigue life due to the presence of fine precipitate in the ferrite matrix coupled with martensite as second phase.

The yield stress of the high strength dual phase steel sheet obtained is 350-500 MPa. The tensile strength obtained is min. 600 MPa. The min. uniform elongation is 16% and 22% minimum total elongation.

Further, strain hardening exponent (“n”) of the high strength dual phase steel sheet is 0.15-0.16. Yield to Tensile strength (ratio) of the dual phase steel is 0.6-0.8 and the hole expansion ratio in punched condition is about 40%.

Experimental Analysis

For the purpose of example only, a slab of the composition (given in Table 1) according to the process 100 (Strip 1) was continuously cast in a CSP mill. Slab was hot rolled. The ROT cooling was done in accordance with the present disclosure and the cooling profile is given in FIG. 2. The mechanical properties steels sheet are listed in Table 2, 3 & 4. The microstructures of the steels are shown in FIGS. 4, 5, 6 & 7. It is clear from the mechanical properties and the microstructures achieved, that the target properties can be achieved when the chemistry and ROT cooling parameters do conform to the requirements of the disclosure.

The optical (both Nital and Le pera etched) and SEM microstructures are presented in FIGS. 4, 5, 6 & 7 which consist of ferrite and martensite. Tensile test samples with 50 mm gauge length were prepared in accordance to ASTM E8 standard. Typical tensile test plot is given in FIG. 3. It is evident from the figure and table that newly developed steel has minimum 600 MPa tensile strength, 16% uniform elongation and minimum 22% total elongation, the strip has high strain hardening co-efficient 0.15, yield ratio (Yield strength to Tensile strength) between 0.6 & 0.8. The steel has dispersion of fine precipitates in ferrite matrix. The identities of these precipitates are confirmed using Energy Dispersive Spectroscopy (EDS) and Selective Area Diffraction (SAD) techniques in TEM. The precipitates are majorly Nb(C,N) as described in FIG. 8a-f . The steel also has very fine average grain size below 3 μm.

TABLE 1 Composition (in wt. %) of the samples tested C Mn Si Cr Nb Al P S N Strip 0.058 1.18 0.02 0.59 0.02 0.036 0.02 0.003 <0.005 1

TABLE 2 Tensile properties of the tested samples TS TEL Sample ID Description YS (MPa) (MPa) (%) YS/TS N Strip 1 Rolling 441 652 25 0.68 0.16 Direction (RD) 457 666 25 0.69 0.16 437 653 24 0.67 0.16 Average RD 445 657 25 0.68 0.16 Transverse 447 658 22 0.68 0.15 Direction 458 655 23 0.70 0.15 (TD) 444 655 22 0.68 0.16 Average TD 450 656 22 0.69 0.15

TABLE 3 Hole expansion of the tested samples HER = Sample Punched Expanded (D2 − Sample ID No. Dia. (D1) Dia (D2) D1)/D1 HER * 100 (%) Remarks Strip 1 1 10.04 13.48 0.34 34.0 multiple expansion 2 10.04 14.85 0.48 48.3 too large crack 3 10.03 14.88 0.49 49.2 too large crack 4 10.04 14.14 0.40 40.0 Single small crack 5 10.03 13.97 0.39 38.5 Single small crack

TABLE 4 Quantification of microstructural constituents Hardness Ferrite Martensite Bainite Grain size (VHN with Sample ID (%) (%) (%) (μm) 0.1 kgf) Strip 1 80 17 <5 2-3 208 ± 10 

1. A process for producing dual phase steel sheet, comprising steps of: making a liquid steel having chemical composition in wt % of C: 0.03-0.12, Mn: 0.8-1.5, Si:<0.1, Cr: 0.3-0.7, S: 0,008 maximum, P: 0.025 maximum, Al: 0.01 to 0.1, N: 0.007 maximum, Nb: 0.005-0.035, and V: 0.06 maximum remainder Fe and inevitable impurities; continuous casting the liquid steel into a slab; hot rolling the slab into a hot rolled sheet at finish rolling temperature (FRT) of 810-870° C. cooling the hot rolled sheet on the run out table at a cooling rate 40-70° C./s to an intermediate temperature (TINT) of 720° C.≤TINT≤650° C.; natural cooling the hot rolled sheet for a duration of 5-7 seconds; and rapidly cooling the hot rolled sheet to transform remaining carbon enriched austenite to martensite, at cooling rate of 40-70° C./s to a coiling temperature below 400° C. to produce a dual phase steel sheet.
 2. The process as claimed in claim 1, wherein the slab is re-heated to a temperature of 1100-1200° C. to dissolve precipitates before hot rolling.
 3. The process as claimed in claim 1, wherein a yield stress of the dual phase steel sheet is 350-500 MPa.
 4. The process as claimed in claim 1, wherein the dual phase steel sheet has a 600 MPa minimum tensile strength.
 5. The process as claimed in claim 1, wherein the dual phase steel sheet has 16% minimum uniform elongation.
 6. The process as claimed in claim 1, wherein the dual phase steel sheet has 22% minimum total elongation.
 7. The process as claimed in claim 1, wherein a strain hardening exponent (n) of the dual phase steel sheet is 0.15-0.16
 8. The process as claimed in claim 1, wherein a yield strength to tensile strength ratio of the dual phase steel is 0.6-0.8.
 9. The process as claimed in claim 1, wherein the dual phase steel sheet has a hole expansion ratio in a punched condition is of about 40%.
 10. The process as claimed in claim 1, wherein the dual phase steel sheet 75-90% ferrite, 10-25% martensite and <5% bainite by volume.
 11. The process as claimed in claim 1, wherein a grain size of the dual phase steel sheet is 2-5 μm.
 12. A dual phase steel sheet, comprising: a chemical composition in wt % C: 0.03-0.12, Mn: 0.8-1.5, Si:<0.1, Cr: 0.3-0.7, S: 0,008 maximum, P: 0.025 maximum, Al: 0.01 to 0.1, N: 0.007 maximum, Nb: 0.005-0.035, and V: 0.06 maximum remainder Fe and inevitable impurities;
 13. The dual phase steel sheet as claimed in claim 12, wherein a yield stress of the dual phase steel sheet is 350-500 MPa.
 14. The dual phase steel sheet as claimed in claim 12, wherein the dual phase steel sheet has a 600 MPa minimum tensile strength.
 15. The dual phase steel sheet as claimed in claim 12, wherein the dual phase steel sheet has 16% minimum uniform elongation.
 16. The dual phase steel sheet as claimed in claim 12, wherein the dual phase steel sheet has 22% minimum total elongation.
 17. The dual phase steel sheet as claimed in claim 12, wherein a strain hardening exponent (n) of the dual phase steel sheet is 0.15-0.16.
 18. The dual phase steel sheet as claimed in claim 12, wherein a yield strength to tensile strength ratio of the dual phase steel sheet 0.6-0.8.
 19. The dual phase steel sheet as claimed in claim 12, wherein a hole expansion ratio in a punched condition is about 40%.
 20. The dual phase steel sheet as claimed in claim 12, wherein the dual phase steel sheet has 75-90% ferrite, 10-25% martensite and <5% bainite by volume.
 21. The dual phase steel sheet as claimed in claim 12, wherein a grain size of the dual phase steel sheet is 2-5 μm. 